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Materials Science and Engineer

Electrochemical formation of highly pitting resistant passive films on abiomedical grade 316LVM stainless steel surface

Abdullah Shahryari a, Sasha Omanovic a,⁎, Jerzy A. Szpunar b

a Department of Chemical Engineering, McGill University, Montreal, Quebec, Canada H3A 2B2b Department of Mining and Materials, McGill University, Montreal, Quebec, Canada H3A 2B2

Received 13 July 2007; received in revised form 31 August 2007; accepted 12 September 2007Available online 20 September 2007

Abstract

The results discussed in the paper demonstrate that a significant improvement in pitting corrosion resistance of a biomedical grade 316LVMstainless steel can be achieved by electrochemically forming highly-protective passive oxide films on the material's surface, under cyclicpotentiodynamic polarization conditions. The film formed in a sodium nitrate electrolyte is completely resistant to pitting corrosion in simulatingphysiological solutions even at high temperatures (60 °C), and after sterilization. The high pitting resistance of the electrochemically-formed filmswas explained on the basis of their semiconducting properties. Namely, the enrichment of the outer part of the electrochemically formed passivefilm with Cr(VI)-species results in a decrease in the density of oxygen vacancies, which act as pitting initiation sites, and their ‘replacement’ bymetal vacancies formed by the electrochemical oxidation of Cr(III) to Cr(VI). In this configuration, the outer Cr(VI)-rich oxide layer behaves ascation selective, which results in the increased pitting corrosion resistance of the film. The simple electrochemical passivation technique discussedin the paper can be efficiently used to form highly pitting resistant passive films on 316LVM-built medical implant devices of any geometry.© 2007 Elsevier B.V. All rights reserved.

Keywords: Biomaterials; Stainless steel; Pitting corrosion; Passive films; Electrochemical potentiodynamic passivation; Semiconducting properties

1. Introduction

Stainless steels (SSs) are widely used as biomaterials andmaterials of construction. In biomedical applications they areused as coronary and pulmonary stents, hip prosthesis, screws,external fixations, etc. This is mainly due to their goodresistance to general corrosion. Generally, the first requirementof any material serving in a biological system is that it should beinert and not cause any undesirable reaction with its surround-ing. When SS is placed inside a tissue, the interaction betweenthe implant and the tissue determines the degree of itsbiocompatibility. Corrosion, as an electrochemical process,commences on the surface of SS implants and subsequentlyaffects the body's response. This is undesirably followed byrelease of ions such as chromium and nickel in the surrounding

⁎ Corresponding author. Department of Chemical Engineering, McGillUniversity, 3610 University Street, Montreal, QC, Canada H3A 2B2. Tel.: +1514 398 4273; fax: +1 514 398 6678.

E-mail address: sasha.omanovic@mcgill.ca (S. Omanovic).

0928-4931/$ - see front matter © 2007 Elsevier B.V. All rights reserved.doi:10.1016/j.msec.2007.09.002

tissue which can negatively affect the response of the host bodyto the biomaterial. Okazaki et al. [1] studied the degradation rateof a number of commonly used biomaterials in differentbiological solutions. In the case of SS316LVM, they reported arelatively high release rate of the alloying elements, mainlychromium and iron, into the solution, reaching a concentrationof 2 μg cm−2 after a seven-day exposure. The interaction ofvarious elements, originated from dissolution of SS, with thehuman body has attracted much attention and has beenextensively studied [2–4]. The literature has proven that anydeviation from the natural concentration of the ions such aschromium, iron and nickel inside the body may severely preventthe body from functioning properly, and can ultimately lead tosevere health-related consequences. Therefore, corrosion of thematerial is the first issue to be considered when it is designed tofunction in a human body.

Although, the resistance of SS passive films to generalcorrosion is relatively high, the films are highly susceptible tolocalized forms of corrosion. Pitting corrosion is one of the mostsevere types of localized attack on SSs, which can limit their

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application in bio-systems. Pitting corrosion can adversely affectboth biocompatibility and mechanical strength of the implant.For instance, it has been shown that pitting corrosion initiated onthe surface of a coronary stent strut can lead to a completemechanical failure of the stent [5]. Shih et al. [6] haveinvestigated the influence of the surface passive oxide filmproperties on the pitting resistance of SS stents. They showed byin vivo experiments that the pitting resistance of the implantedstents in the dog's abdominal aorta is directly dependent on thestent material's passive film properties. Further, the thrombo-genic activity of the SS surface is also believed to be dependenton the biomaterial's passive oxide film physico-chemicalproperties and its pitting resistance [7–9]. Investigation of thethrombogenicity of variously treated SS cardiovascular stentshas shown that the minimum thrombosis is detected on theimplanted stent exhibiting the highest resistance to pittingcorrosion in in vivo experiments [7–9]. Further, SS fracturefixation parts such as screw heads have been frequently reportedto fail due to pitting corrosion [10,11]. Also, localized forms ofcorrosion on hip prostheses are one of the commonly reporteddefects on the retrieved parts [12]. One of the other majorproblems related to pitting corrosion of orthopedic implants isthe elevated friction originated from the pitting corrosionproducts acting as debris at the joints. Release of corrosionproducts into the tissue surrounding the implant can elicitdifferent types of reactions in the host tissue. The increase of thelocal and overall concentration of certain species associated withcorrosion of the implants has been cited in many clinical repots[13–15]. In the case of SSs, Williams et al. [16] showed that atthe screw-plate junctions of the bone fixation components, signsof inflammations along with the corrosion products, iron-containing granules and micro-plates containing chromiumcompounds, have been found.

The resistance of SS to corrosion is owed to the presence of athin oxide film on its surface, known as passive film. Thephysico-chemical properties of this passive film control thematerial's corrosion behavior, its interaction with the body, andthus the degree of the material's biocompatibility. A number ofdifferent properties of the passive film, such as chemicalcomposition, electronic properties, thickness, etc. determine thecapability of the SS to resist pitting. It is a common agreementthat the passive film on a biomedical grade SS, e.g. 316LVM,has a duplex structure, being composed of Cr-oxide and Fe-oxide sub-layers. However, the passive/protective nature of thepassive film is generally attributed to the presence of Cr-oxide.Consequently, it is believed that an increased content ofchromium in the passive film results in an increased resistanceof the SS to pitting corrosion [17–20]. However, Shahryari et al.[21] have shown that the enrichment of electrochemicallyformed passive films with Cr(III) does not yield an improve-ment in their pitting resistance relative to the naturally grownpassive film, but rather their the enrichment with Cr(VI) species.It has also been frequently reported that molybdenum plays apositive role in pitting corrosion [22,23]. Merello et al. [24]claimed that presence of Mo in SS in the amount higher than1 wt.% results in a considerable improvement in their pittingresistance of the material. On the other hand, the presence of

some inclusions on the SS surface, such as sulfide inclusions,represents a major problem due to their acceleration of pittingcorrosion [25–27]. Ryan et al. [26] suggested that the depletionof chromium around theMnS inclusions triggers the initial stageof pitting corrosion. This has been questioned by Meng et al.[28] and Schmuki et al. [29], who suggested that no chromium-depleted zone was detected around any of the MnS inclusions.While it is known that low-sulphur steels (e.g. a biomedicalgrade 316LVM) have improved corrosion resistance, someresearchers believe that pitting corrosion in these steels is alsoalways associated with MnS inclusions [26].

Significant efforts have been made to develop methodsfor modification of the SS surface and/or its passive films inorder to improve the material's pitting corrosion resistance, andthus its biocompatibility. These methods have been focusedmainly on the removal of surface inclusions [17,30], on themodification of chemical properties and element distribution inthe passive film [31–33], or on the increase in the Cr/Fe ratio inthe film [17–20]. Some improvement in pitting resistance hasbeen achieved by these and some other methods [17,19,34,35].Nitric acid has been used as one of the most popular chemicalpassivation reagents for surface treatment of surgical SSimplants [36]. The literature has emphasized a beneficial effectof nitric acid on chromium enrichment in the modified passivelayer [17,19,30]. Coating of 316 L SS with hydroxyapatite hasalso been used in order to increase the material's biocompat-ibility in terms of general and pitting corrosion resistance [37].Shih et al. [6,18] have tested the properties of a 316 L surfacemodified using various passivation methods, and have con-cluded that the method that produces an amorphous surfaceoxide layer gives the highest increase in the material'sbiocompatibility. Laser surface modification has been studiedby many researchers as a method of improving the pittingresistance of the SS surface [33,38–41]. According to theliterature, the general idea on the influence of laser surfacetreatment is that this method produces a homogeneous surfacehaving a fine-grained structure, and also dissolves or confinesthe carbide particles. Electrochemical polishing of 316 L SSslotted tube coronary stents has been used as a surfacepretreatment method in order to increase the material'sbiocompatibility [42]. Nitriding, a surface treatment methodcommonly used to increase the wear resistance of SSs, has alsobeen used as a SS surface modification method [43,44]. Ionimplantation of the 304 SS surface by molybdenum, yttrium,titanium and nitrogen has also been applied to increase thematerial's corrosion resistance [45–47].

The goal of the research presented in this paper has been toinvestigate the suitability of a simple electrochemical technique,cyclic voltammetry, for the formation of a highly pitting resistantpassive film on a biomedical grade 316LVM SS surface. Theinfluence of a range of experimental parameters (temperature,chloride concentration, number of polarization scans, etc.) onthe resulting pitting behavior of the material is discussed. It willbe shown that the formed passive film offers a significantlyhigher pitting corrosion resistance than the naturally grownpassive film, thus offering an increased biocompatibility ofthe material. The latter has also been confirmed by preliminary

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experiments (not shown) in which a 95% increase in surfacedensity of pre-osteoblast cells (MC3T3) has been achieved onthe electrochemically modified 316LVM surface. The improve-ment in pitting corrosion resistance of the modified surface hasbeen explained on the basis of chemical composition andsemiconducting properties of the film.

2. Experimental

All the electrochemical measurements were carried out in asingle-compartment electrochemical cell. The counter electrode(CE) was a platinum wire, and a saturated calomel electrode(SCE) was used as the reference electrode (RE). Stainless steelSS316LVM was used as the working electrode (WE). Itschemical composition is shown in Table 1. The WE wasprepared by embedding SS316LVM pieces into epoxy resin,exposing a two-dimensional rectangular cross-section of0.36 cm2 to the electrolyte.

Several different electrolytes were used in this research. Forpassivation of 316LVM, 0.1 M NaNO3 (NN) and 0.1 Mphosphate buffer pH 7.4 (PB) solutions were used as passivatingelectrolytes. The latter was also used in capacitance measure-ments. Three different corrosion testing electrolytes were used:pure saline solution (0.16 M NaCl), phosphate buffer saline(PBS), and Hank's solution. All the solutions were preparedusing deionized water having resistivity of 18.2 MΩ cm.

All the electrochemical testing experiments were performedin oxygen-free electrolytes which was achieved by continuouslypurging the electrolyte with argon, starting 30 min prior tothe measurement and continuing during the measurement. Priorto each experiment, the working electrode surface was wetpolished using 600 grit abrasive sandpaper followed bysonication in deionized water to remove any polishing residues.The specimens were next degreased with ethanol and immersedin the cell containing an appropriate electrolyte. This wasfollowed either by leaving the WE at open-circuit potential(OCP) for 1 h before the subsequent measurements (this surfaceis named here as “unmodified”), or by electrochemicallypassivating the 316LVM surface prior to the OCP measurementin either 0.1 M NaNO3 or 0.1 M phosphate buffer pH 7.4 (thesesamples are named here as “ME-NN” or “ME-PB”, respective-ly). The electrochemical passivation was done by polarizing theWE between two potential limits specified in the paper for eachmeasurement and at a sweep rate of 100 mV s−1. Prior to eachpassivation experiment, in order to reduce the surface oxidesformed spontaneously during the 316LVM surface polishingand degreasing, the WE was cathodically polarized at −1.0 Vfor 5 min.

The first step in characterization of the surface corrosionresistance was the electrode stabilization at OCP in a corrosion

Table 1Chemical composition of AISI 316LVM stainless steel (wt.%)

Fe Cr Ni C Mo Mn S Si

Bal 16.57 10.34 0.016 2.13 1.54 0.001 0.54

testing solution for a period of 1 h. Next, pitting experimentswere done by anodically polarizing the working electrode from50 mV negative of the OCP to the potential at which a currentdensity of 1 mA cm−2 was reached. This was immediatelyfollowed by a reverse cathodic polarization bias to the startingpotential. Chronoamperometric measurements were performedfollowing the electrode stabilization at OCP in a corrosiontesting solution by potentiostatically polarizing the WE andmeasuring the resulting current. To investigate the semicon-ducting properties of naturally grown and electrochemicallyformed passive films, capacitance measurements were carriedout in the PB solution under a potentiostatic control andcathodic (negative) bias, starting from 0.9 V down to −0.9 V.The applied ac amplitude was ±10 mV.

X-ray photoelectron spectroscopy (XPS) measurementswere made using a VG instrument Escalab 220i XL equippedwith an argon ion gun. The X-ray unmonochromatic source wasAl (1486.6 eV). The ion etching beam was used at 3 keV with amagnification of ten, which provided an etching area of1.5×1.5 mm2. The etching rate was 2 nm min−1 and thepressure during the etching was kept at 10−8 mbar. Thereference energies used for calibration of the binding energieswere the Ag3d5/2 signal at 367.9 eV and the Cu2p3/2 signal at932.7 eV. The analyzer was fixed at normal position (90°) to thesurface. A survey spectrum was first recorded to identify allelements present on the sample surface, followed by recordinghigh resolution spectra. The spectra were fitted using linearbackground subtraction and a combination of Gaussian andLorentzian line shapes with addition of an asymmetry factor forthe metal peaks. In order to derive a quantitative analysis, 2pspectra of the selected elements were recorded. Next, thecalculated area under the deconvoluted peaks after backgroundsubtraction was correlated to the atomic concentration of thecorresponding element using the related correction factors.

3. Results and discussion

3.1. Cyclic polarization

Fig. 1 shows the cyclic voltammogram (CV) of a freshlypolished 316LVM SS electrode recorded in a potential regionbetween −0.8 V and 0.9 V, in which the solid curve representsthe 1st polarization sweep and the dashed curve represents the200th polarization sweep. The shape of the CV changes greatlywith the number of polarization sweeps. The CV recorded in thefirst sweep (solid line) displays an anodic hump (A1) at ca.−0.5 Vand a broad current shoulder (A2) cantered at ca. 0.5 V. Inthe reverse cathodic polarization for the 1st sweep, thevoltammogram shows two reduction peaks, C1 and C2. Therelation of the shoulders/peaks to specific redox reactions has

P Cu Sn Co N O Nb

0.024 0.28 0.009 0.09 0.03 34 ppm 0.01

Fig. 2. Pitting polarization curves of the unmodified, ME-NN and ME-PB316LVM surfaces recorded at 22 °C in a 0.16 M PBS solution pH 7.4. Scanrate=1 mV s−1. Modification of the surfaces was done by potentiodynamiccyclic polarization of the electrode at 22 °C in 0.1 M sodium nitrate (ME-NN) orphosphate buffer (ME-PB) between −0.8 Vand 0.9 Vat a scan rate of 100 mV s−1

by applying 300 sweeps.

Fig. 1. Cyclic voltammograms of the 316LVM surface recorded in 0.1 MNaNO3. The solid curve shows the 1st polarization sweep and the dashed curveshows the 200th polarization sweep. Scan rate=100 mV s−1. Inset: Variationof the anodic-to-cathodic total charge ratio with the number of sweeps.Temperature=22 °C.

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already been well explained in the literature [48–51]. Briefly,the anodic peak (A1) is related to the formation of a Fe(II)/Fe(III) layer (i.e. Fe2O3) on the pre-existing Cr(III)-oxide surfacelayer, followed by the oxidation of Cr(III)-oxide to soluble Cr(VI)-species (A2), most likely Cr2O7

2−. The cathodic peak (C2) isrelated to the reduction of Cr(VI)-species back to Cr(III)-oxide,and (C1) to the reductive decomposition of an oxide layercomposed of Fe2O3 back to Fe(II)-species.

On the freshly polished 316LVM surface (1st sweep, solidcurve), anodically formed Cr(VI)-species are mainly lost insolution by diffusion through a very thin and non-compact pre-formed Fe2O3 passive film, with subsequent dissolution in theaqueous phase. This explains the absence of passive transitionassociated with shoulder A2 in the anodic scan, and poordefinition (small change) of the cathodic peak C2 (solid curve)in the cathodic scan. On the other hand, the 200th sweep in Fig. 1(dashed line) shows that the charge associated with the Cr(III)-to-Cr(VI) transition in the potential region of A2 and C2 isnegligible compared to that recorded in the first sweep. Thisdemonstrates that Cr(VI) species formed in the 200th anodicsweep remain ‘arrested’ in the growing surface passive film. Italso indicates that the passive film formed during the prolongedcyclization of the electrode is more compact and thicker than thefilm formed in the 1st sweep, thus effectively preventing thedissolution of Cr(VI) species into the solution. As it will beshown later in the text (Fig. 10), the major product of chromiumoxidation during the electrode cyclization is Cr(III), formed byoxidation of metallic chromium, with a small contribution of Cr(VI) species accumulated in the outer part of the passive film.

In order to get a better quantitative insight into changesduring the cyclization of the 316LVM electrode, the ratio ofthe total anodic-to-cathodic charge is presented as an inset toFig. 1. The curve shows that the anodic reactions related toshoulders A1 and A2 appearing in the 1st sweep are quiteirreversible (QA /QC=11.09) and that only ca. 9% of the chargerelated to these anodic processes is used to form species that donot dissolve into the solution but form the surface passive oxide

film. However, by cyclization of the surface, the reversibility ofthe anodic processes rapidly increases, and after ca. 20 sweepsit approaches almost 100% (QA /QC=1.05, after the 200thsweep). This indicates that the cyclic polarization of the316LVM surface under the given conditions significantlycontributes to the enhancement of the passive properties ofthe surface oxide film, preventing the dissolution of theunderlying metals into the solution even at high anodicpotentials. Hence, this could be potentially used as an efficientand convenient method for the formation of passive films on the316LVM surface that offer better corrosion resistance thannaturally grown passive films. In order to verify this hypothesis,pitting corrosion polarization measurements have been madeon both the unmodified surface (on which the passive filmwas grown naturally) and on surfaces modified under variouspassivation conditions.

3.2. Pitting polarization

Measurements of pitting resistance of modified and unmod-ified surfaces were performed in various solutions: pure saline,PBS, and Hank's. Fig. 2 shows the anodic polarization curves ofthe 316LVM surface modified in two different electrolytes(sodium nitrate, ME-NN, and phosphate buffer, ME-PB) andalso the response of the unmodified surface. The unmodifiedsurface represents the 316LVM surface on which the passivefilm was grown naturally during the stabilization of theelectrode at OCP. All the curves were recorded in PBS. Thepitting resistance of the unmodified surface is in agreementwith the literature [26,52–54] with the onset of pitting at ca.0.4 V. On the other hand, the polarization curve recorded onthe ME-PB surface shows a considerable improvement inpitting resistance. The passive region extends to high anodicpotentials, and the breakdown of the passive film commencesat ca. 1.05 V. The examination of the ME-PB electrode after the

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polarization measurement revealed that the increase in currentobserved at ca. 1.05 V is actually not due to pitting but due tothe crevice corrosion at the electrode/resin interface. Thesurprising result was that the polarization curve for the ME-NNsurface shows the absence of a pitting-loop, which demonstratesthat the surface passive film formed in NN is completelyresistant to pitting under the experimental conditions applied.The rise in current at ca. 1.0 V is not due to pitting, but due tooxygen evolution. In addition to the high pitting resistance, theother positive effects observed on the ME-NN surface are adecrease in passive current by ca. 1.5 orders of magnitude, anda shift in OCP to more noble values by ca. 160 mV. Theexperiments performed in Hank's solution also demonstratedthe absence of pitting on the ME-NN surface, while the surfaceon which the passive film was grown naturally (unmodifiedsurface) pitted already at 0.16 V (not shown). The sameexperiment was also performed with high-sulfur SS 316 L, andthe results demonstrated that the surface on which the passivefilm was grown naturally pitted at 0.4 V, while the surface onwhich the passive film was formed electrochemically in NNoffered a considerably higher pitting corrosion resistance, andpitted only at potentials above 0.96 V (not shown).

The results in Fig. 2 and those in Hank's solutiondemonstrate that the passive film formed electrochemically insodium nitrate (NN) offers higher corrosion resistance than thefilm formed in phosphate buffer (PB). Therefore, further in thetext the ME-NN surface will be termed as the modified surface,unless otherwise stated.

Literature on the modification of SS surfaces by variousmethods report a wide range of controversial results. Forexample, when a 316 L SS surface was thermally modified, noimprovement in pitting corrosion was noted in Ringer's solutionat 37 °C [18]. On the other hand, electropolishing of the samesurface resulted in a shift in pitting potential by about 0.25 V.When the surface was coated with an amorphous oxide, nopitting was observed. However, it should be noted that theauthors in [18] performed these pitting polarization measure-ment at a very high scan rate, 167 mV s−1, which might havebeen too high to induce pitting in the time scale of themeasurement in the pitting-susceptible region. Hence, ouropinion is that these results cannot be taken as completelyconclusive. Further, the authors in [17] showed that passivationof a 316 surface in a nitric acid solution yielded a shift in thepitting potential by about 0.15 V when tested in a neutralaqueous solution containing 1 M of chlorides at 70 °C, while theanodic modification of a 304 SS surface in a metasilicate-containing solution [55] resulted in a positive shift in pittingpotential by about 0.3 V at room temperature. Treatment of thesame surface under the same anodic conditions in a basichydroxide solution showed no improvement in the material'spitting resistance [55]. A significant increase in the pittingpotential (by ca. 0.7 V) was achieved when a 304 L SS surfacewas nitrided, but the drawback was that the resulting passivecurrent was higher by one order of magnitude with respect to theunnitrided surface [44]. Modification of a 304 SS surface by UVradiation improved the pitting resistance of the surface in 0.6 Mchloride solution by about 0.09 V, with a decrease in the pit

generation rate by half an order of magnitude [35]. A slightlyhigher positive displacement in the pitting potential (ca. 0.17 V)was obtained in 0.1 M chloride solution when a 316 SSsurface was anodically polarized at a constant potential and UV-illuminated for 5 h [56]. Using an excimer laser surface treatmentmethod in an air stream, a shift in pitting potential by about 0.15Vwas achieved in 0.6M chloride solution at room temperature for athat of 316LS SS (the chemical compassion of 316LS is close tothat of 316LVM) surface [39]. In addition, a decrease in passivecurrent by about one order of magnitude was achieved. However,the same treatment in a nitrogen atmosphere resulted in asignificant increase in both the pitting corrosion susceptibility andpassive current. On the other hand, the results published in [41]showed that the laser surface melting of 304 SS in a nitrogenatmosphere resulted in an increase in pitting potential by about0.3 V, and by 0.32 V when the surface was treated in an argonatmosphere [40]. This short literature review on the surfacemodification of SSs for the improvement of their resistance topitting corrosion shows that the simple electrochemical modifi-cation procedure presented in this paper offers a very good andconvenient alternative method that could be used to increase theresistance of the 316LVM SS surface to pitting corrosion.

It should be noted that the adsorption of phosphate ions(present in PBS and Hank's solution) on SSs is known tocontribute to corrosion inhibition [57,58]. Hence, the pittingcorrosion resistance of the ME-NN and unmodified surfaceswere also tested in pure saline solutions in the absence ofphosphate ions (discussed in the forthcoming sections). Itshould also be noted that the concentration of chlorides in thepure saline solution is the same as that in PBS and Hank'ssolution (0.16 M), but due to the absence of inhibiting anions,pure saline is a more pitting-corrosion aggressive solution.

3.3. Influence of number of sweeps

The electrochemical passivation process was discussed indetails earlier in the text. It was also explained that thereversibility of the anodic/cathodic surface processes increaseswith an increase in the number of passivation sweeps (Fig. 1).This was then hypothesized to be related to the observedincrease in the surface pitting resistance. In order to verify thishypothesis, the pitting resistance of the surface modified byapplying different number of sweeps was investigated. Fig. 3shows the variation of the pitting potential of 316LVM with thenumber of passivation sweeps applied during modification ofthe surface in the NN solution. There is an obvious correlationbetween the pitting potential recorded in saline solution and theapplied number of sweeps. With an increase in the number ofsweeps, the pitting potential also increases from ca. 0.23 V(naturally grown passive film) to ca. 1.07 V for the passive filmformed by applying 300 sweeps. This is in agreement with thedata presented in the inset to Fig. 1, and indicates that with anincrease in the number of passivation sweeps, the compactnessand thickness of the passive film also increases, thus con-tributing to an increase in pitting resistance of the material. Nofurther improvement was observed by increasing the numberof polarization sweeps above 300. Thus, all the data on the

Fig. 4. Choronoamperometric curves of the (1) unmodified and (2) modified316LVM surface recorded at 22 °C in 0.16 M PBS at 0.35 VSCE (plot a) and0.45 VSCE (plot b) Modification of the surface was done by potentiodynamiccyclic polarization of the electrode at 22 °C in 0.1 M NaNO3 between −0.8 Vand 0.9 V at a scan rate of 100 mV s−1 by applying 300 sweeps.

Fig. 3. Dependence of the pitting potential on the number of sweeps appliedduring the modification of the 316LVM surface. The modification of the surfacewas done by the potentiodynamic cyclic polarization of the electrode at 22 °C in0.1 M NaNO3 between −0.8 V and 0.9 V at a scan rate of 100 mV s−1 byapplying the specified number of sweeps. The corresponding pitting corrosionmeasurements were done in 0.16 M NaCl at a scan rate of 1 mV s−1 and at 22 °C.

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modified surface presented in the remaining sections of thepaper refer to this condition, if not otherwise stated.

3.4. Pitting choronoamperometric measurements

It is a general agreement that pitting corrosion is triggeredfrom defects and/or susceptible spots on the surface [25–28].Therefore, as the number of the active spots on the surfaceincreases, the probability of the stable pits formation alsoincreases. One of the means of comparing the susceptibility ofmaterials to pitting corrosion (in addition to measuring thepitting potential) is to measure the pit initiation frequency, usingchronoamperometry. Generally, potentiostatic polarization of SSin a metastable pitting potential region (between the OCP andpitting potential) results in the formation of metastable pits onthe material's surface, followed by their repassivation, which isin chronoamperometry manifested as a series of sharp currenttransients (spikes) [59]. Fig. 4a shows the chronoamperometrycurve of the unmodified (curve 1) and of the ME-NN (curve 2)316LVM surface recorded at 0.35 V, which is slightly belowthe pitting potential (by ca. 50 mV) of the unmodified surface(Fig. 2). There is an obvious difference between the behavior ofthe unmodified and ME-NN surface. The graph demonstratesthat no current spikes are observed on the curve recorded on theME-NN surface (curve 2). However, large current spikes,representing local dissolution of the passive film followed by itsrepassivation, are visible on the unmodified surface (curve 1).Although the observed current spikes (Fig. 4a, curve 1) seem tobe quite sharp, the actual time length of each spike (from theinitiation to death) is ca. 3 s. This time covers the pit initiationstage, followed by a short pit propagation period. However, at0.35 V (Fig. 4a, curve 1) the propagation (growth) of the pit toform a critical-pit is not thermodynamically and kineticallyfavourable, and the pit eventually repassivates.

For comparison, the results recorded at 0.45 V, which isslightly above the pitting potential (by ca. 50 mV) of theunmodified surface, are presented in Fig. 4b. Again, the formationof metastable pits on the modified surface is completely absent

(curve 2), and the recorded passive current sharply decreasedreaching a quazi-stable value after ca 10 min. Even whenthe surface was polarized for a longer time (1 h), its stabilityremained good and no metastable pits were formed. On the otherhand, curve 1 (Fig. 4b) shows that the current on the unmodifiedsurface first started to decrease due to the initial passivation of thesurface, and then, after approximately 1 min, suddenly started toincrease as a result of the localized dissolution (pitting) of thepassive film. No current spikes could be observed on this curve,indicating that stable pits were formed. A steady-state (localized)dissolution of the film, characterized by high current density (ca.10mA cm−2) was reached already after ca. 15min. The results inFig. 4 are in agreement with those presented in Figs. 2 and 3,and clearly demonstrate that the electrochemically formedpassive film (ME-NN surface) offers a significantly higherpitting corrosion resistance than the naturally grown passive film(unmodified surface).

3.5. Influence of chloride concentration

It was mentioned earlier that the ME-NN modified surfacedid not pit in PBS and in Hank's solution. Also, increasingthe NaCl concentration in those solutions to even 1 M did notresult in pitting of the ME-NN surface. Therefore, in order toevaluate the pitting resistance of the 316LVM surface in moresevere conditions, experiments were performed by testingthe unmodified and ME-NN surfaces in pure saline solutionsof high chloride concentrations (up to 1 M). Fig. 5 showsthe dependence of pitting potential, Epit, of the unmodifiedand ME-NN surfaces on the chloride concentration in thetesting electrolyte. The variation of the pitting potential withchloride concentration shows a similar trend for both surfaces,but there is a large difference in the corresponding absolutepitting potential values. For the unmodified surface, the pittingpotential at a chloride concentration of 0.16 M is 0.23 V. Anincrease in the chloride concentration to 0.5 M and then to 1 Mresults in a decrease in Epit to 0.20 V and 0.08 V, respectively.

Fig. 6. Dependence of pitting potential on the modification solution temperature.The modification of the surface was done by potentiodynamic cyclicpolarization of the electrode in 0.1 M NaNO3 at the specified temperature,between −0.8 Vand 0.9 Vat a scan rate of 100 mV s−1 by applying 300 sweeps.The corresponding pitting measurements were done in 0.16 M NaCl solution at22 °C and at a scan rate of 1 mV s−1.

Fig. 5. Dependence of pitting potential of the ME-NN and unmodified surfaceson the chloride concentration in the NaCl testing solution. Modifications of thesurface was done by potentiodynamic cyclic polarization of the electrode at22 °C in 0.1 M NaNO3 between −0.8 V and 0.9 V at a scan rate of 100 mV s−1

by applying 300 sweeps. The corresponding pitting measurements were done in0.16 M NaCl solution at 22 °C and at a scan rate of 1 mV s−1.

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On the other hand, the modified surface shows a significantlyhigher pitting corrosion resistance, characterized by a very highpitting potential, Epit =1.07 V, which remains constant with anincrease in chloride concentration even to 0.5 M. By furtherincreasing the chloride concentration to very high values, i.e.1 M, the pitting potential decreased to 0.5 V. Nevertheless, itshould be noted that even at such a high concentration ofchlorides in the solution, the modified surface offers asignificantly higher pitting resistance than the unmodifiedsurface, 0.5 V versus 0.08 V, respectively. This shows that theelectrochemical modification of 316LVM results in theformation of a passive film that offers a significantly higherpitting corrosion resistance than the naturally grown passivefilm even at chloride concentrations that are more than six timeshigher than those in the human body. This demonstrates that theapplied electrochemical passivation procedure could also beused to increase the pitting corrosion resistance of 316LVM foruse in some other applications such as marine, industrial, etc.

3.6. Influence of temperature

The importance of temperature on the SS pitting corrosionresistance has been frequently reported in the literature [59–61].It has been shown that an increase in temperature results in asignificantly higher frequency of metastable pit formation[49,62,63]. An increased pitting corrosion susceptibility of SSat higher temperatures can considerably lower the material'ssuitability in biomedical applications, especially when used asan implant material. Therefore, it is important to investigate thetemperature-dependent pitting behavior of 316LMV studied inthis work.

Fig. 6 demonstrates the influence of the modificationsolution temperature on the resulting pitting corrosion behaviorof the electrochemically formed passive film (ME-NN). Thepitting measurements were performed at 22 °C in 0.16 M NaCl.The graph shows that an increase in the modification solutiontemperature from 22 °C to 60 °C resulted only in a slightdecrease in the pitting potential, 1.07 V, 1.05 V and 0.95 V,

respectively. This small decrease in the pitting potential with anincrease in the film formation temperature could be due to thedifference between the passive film formation temperature andtesting temperature (38 °C). Namely, with an increase in thefilm formation temperature the specific volume of the formedpassive film also increases (the films expands). Then, after thesample is cooled down to the testing solution temperature(23 °C), the passive film shrinks. This can result in theformation of nano-sized cracks (pores) in the film, whichrepresent pitting-susceptible surface sites. With an increase inthe film formation temperature, the surface density of thesecracks also increases and consequently, the pitting potentialdecreases. However, Fig. 6 demonstrates that this phenomenonis almost marginal in the temperature range investigated, i.e. itdoes not have a significant influence on the resulting pittingresistance of the surface passive film. Nevertheless, Fig. 6shows that even the passive film formed at 60 °C offers asignificantly higher pitting potential (0.95 V) than the naturallygrown passive film (0.22 V). Pallotta et al. [49] and Cristofaroet al. [64] have shown that the formation of a passive layer onSS at higher temperatures promotes the chance of replacementof Cr(III) by Fe(II). They have postulated that this, in turn,results in a decrease in the Cr/Fe ratio in the film and,consequently, in a decrease in the pitting resistance and aconsiderable increase in the passive current of the material(0.55 VSCE and 0.3 VSCE at 15 °C and 60 °C, respectively).However, as we will demonstrate later in the text, the increasedpitting resistance of the electrochemically modified 316LVMsurface does not originate from the increased Cr/Fe ratio in thepassive film, which thus explains the difference between ourresults (Fig. 6) and those in [49,64].

The previous section describes the influence of the modi-fication solution temperature on the corresponding pittingresistance of the material, which we have used to optimize thepassive film formation conditions. However, it would also beinteresting to investigate the influence of the corrosive solutiontemperature on the material's pitting resistance, especially inthe region of physiological importance (36–42 °C). This is of a

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practical importance for the use of 316LVM as an implantmaterial due to possible temperature changes occurring in theimplanted material's surrounding as the consequence of inflam-mation. Fig. 7 shows the dependence of pitting potential on thetesting solution temperature. The corresponding passive film wasformed at 22 °C. The graph shows that an increase in the testingsolution temperature resulted in a decrease in the pitting potentialfor the electrochemically formed passive film (ME-NN). Asimilar trend was also recorded for the naturally grown passivefilm (unmodified), although this is not clearly visible on the graphdue to the difference in the pitting potential values recorded on thetwo surfaces. The enhanced susceptibility of SSs toward pitting athigher temperatures has been reported in the literature [59,61,65].The trend observed in Fig. 7 is in accordance with an Arrhenius-like type of kinetics, i.e.with an increase in temperature the rate ofthe corrosion process also increases, resulting in a decreasedresistance of the material towards pitting corrosion. Nevertheless,even at 60 °C the electrochemically formed passive film (ME-NN) still offers higher pitting resistance than the unmodifiedsurface, Fig. 7. It is important to mention that when the sameexperiments were performed in Hank's solution, no pitting wasrecorded in the whole temperature region studied. Hence, in thephysiologically important temperature range (36–42 °C) themodified surface is completely resistant to pitting corrosion whentested in a physiological electrolyte (Hank's).

In order to determine the influence of sterilization of the ME-NN modified 316LVM surface on its pitting resistance, thecorresponding samples were sterilized in an autoclave and 70%ethanol. In both cases the film maintained its high corrosionresistance, and pitted at very high anodic potentials, 1.11 Vand1.37 V, respectively (not shown).

The results presented so far demonstrate that the cyclicpotentiodynamic polarization of the 316LVM surface resultsin the formation of a surface passive film that provides aconsiderably higher resistance to pitting than the naturallygrown film. The film maintains its high pitting resistance at high

Fig. 7. Dependence of pitting potential on the testing solution temperature. Thecorresponding pitting measurements were done in 0.16 M NaCl solution at ascan rate of 1 mV s−1 and at the specified temperature. The modification of thesurface was done by potentiodynamic cyclic polarization of the electrode at22 °C in 0.1 M NaNO3 between −0.8 V and 0.9 V at a scan rate of 100 mV s−1

by applying 300 sweeps. The passive film on the unmodified surface was grownnaturally at OCP by keeping the surface in 0.1 M NaNO3 at 22 °C for the samelength of time as the ME-NN surface.

testing solution temperatures and after sterilization. It would benow interesting to determine the major factors that areresponsible for the observed increase in pitting resistance ofthe material. Possible reasons for the very high pitting corrosionresistance of the electrochemically formed passive film could berelated to either a change in the semiconducting properties ofthe passive film, and/or an increase in the film thickness, and/or,as commonly believed [17–20], enrichment of the passive filmwith chromium. In order to investigate which effects areresponsible for the observed high pitting corrosion resistance,XPS and Mott–Schottky measurements were made.

3.7. XPS results

The chemical composition of the passive films formedelectrochemically in sodium nitrate (ME-NN) and phosphatebuffer (ME-PB), and that of the naturally grown passive film(unmodified) was analyzed by XPS.

Fig. 8 shows an example of deconvoluted XPS spectra of theouter part of the passive film for the (a) unmodified and (b) ME-NN sample (chromium and iron responses). The agreementbetween the modeled and experimental data is very good. Thecomparison of the corresponding spectra reveals that Cr(VI)species were detected only in the electrochemically formedpassive film, while Fe(0) species were detected only in thenaturally grown passive film. The predominant state ofchromium in both passive films is Cr(III) in Cr(OH)3, with asmall contribution of Cr(III) in Cr2O3. Both passive filmscontain Fe(II) and Fe(III) mostly in Fe3O4, with somecontribution of Fe(II) in FeO. The peaks at 715.3, 713.4 and717.2 eVare well-known shake-up satellite peaks [66–69]. Theshake-up satellite is caused by an incident X-ray photontransferring a discrete portion of its energy to the excitation of asecond electron rather than imparting its entire quantum ofenergy to the primary, photo-ejected electron. This photoelec-tron looses a small amount of energy and appears on the XPSspectrum at a slightly higher binding energy [66]. Iron-oxidesatellite structures are frequently used as fingerprints to identifyiron-oxide phases. The shake-up satellite in Fig. 8 at 715.3 eVcan be assigned to Fe(II) in FeO [66,67], while the other twosatellite peaks (713.4 and 717.2 eV) are attributed to Fe(II) inchromite (FeCr2O7) [70]. No shake-up satellite structures arevisible on the chromium spectra. The modeled XPS spectrawere further analyzed to evaluate the influence of the depthdistribution of various Cr and Fe species on the correspondingpitting behavior of the passive film.

Fig. 9(a) shows the atomic Cr/Fe ratio in the passive film forthe unmodified surface and for the surfaces modified in twodifferent electrolytes (ME-NN and ME-PB). For each case, anincrease in the passive film's Cr/Fe ratio with respect to the bulkmaterial is recorded. The highest relative increase in Cr wasobtained in the ME-PB passive film, while the Cr/Fe ratio in theME-NN passive film was even lower than that in the naturallygrown film (unmodified). Nevertheless, the ME-NN surfaceoffered the highest pitting resistance, with a complete absenceof pitting in PBS and Hank's solution. The ME-PB surface alsooffered high resistance to pitting, but unlike the ME-NN

Fig. 8. Examples of deconvoluted XPS spectra recorded on (a) the 316LVM surface on which the passive film was grown naturally (unmodified), and on (b) the316LVM surface on which the passive film was formed electrochemically at 22 °C by potentiodynamic cyclic polarization of the electrode in 0.1 M NaNO3 between−0.8 Vand 0.9 Vat a scan rate of 100 mV s−1 by applying 300 sweeps. The response was recorded after removing the outer passive oxide layer by sputtering for 1 min.All the components in the graphs are labeled based on binding energies of the peaks.

102 A. Shahryari et al. / Materials Science and Engineering C 28 (2008) 94–106

surface, it eventually pitted in PBS when the polarization curvereached high anodic potentials, ca. 1.1 V (Fig. 2). This clearlydemonstrates that the Cr/Fe ratio in the passive film is notthe major factor responsible for the high pitting resistance ofthe electrochemically formed passive films, as usually thought[17–20].

Fig. 9(b) shows the oxygen profile in the electrochemicallyformed (ME-NN and ME-PB) and naturally grown (unmodi-fied) passive film. The oxygen content in all three passive filmsincreases going from the metal/oxide interface toward the outerpassive film surface. This trend is in accordance with thedistribution of oxidized Cr and Fe species in the film (i.e. Cr(III+VI) and Fe(II+III)). Namely, the total amount of oxidized Feand Cr species in the film increases going from the metal/oxideto the outer oxide passive film surface, while the amount of

metallic Fe and Cr decreases and reaches zero at the filmthickness of ca. 12 nm and 6 nm for the two electrochemicallyformed films and the naturally grown film, respectively. Thiscould be seen in Fig. 10 for chromium, and a similar trend wasalso obtained for iron (not shown). Further, a significantdifference between the maximum oxygen content in theelectrochemically formed passive films (ME-NN and ME-PB)and the naturally grown passive film (unmodified) is evident,Fig. 9(b). One of the reasons for this is the formation of Cr(VI)species in the form of Cr2O7

2− in the two electrochemicallyformed passive films (Fig. 10). This species yields a higheroxygen/Cr ratio than Cr(OH)3 and Cr2O3 species that constitutethe naturally grown film (Figs. 8 and 10). Further, the normalizedoxygen content (i.e. the total amount of oxygen in the passivefilm divided by the total film thickness) in the ME-PB and ME-

Fig. 9. (a) Cr/Fe atomic ratio, and (b) oxygen depth profile in the 316LVMpassive film grown naturally (unmodified), and formed electrochemically at22 °C by potentiodynamic cyclic polarization of the electrode in 0.1 M NaNO3

(ME-NN) and 0.1 M phosphate buffer pH 7.4 (ME-PB) between −0.8 V and0.9 V at a scan rate of 100 mV s−1 by applying 300 sweeps.

Fig. 10. Depth profile of different oxidation states of chromium in the passivefilms formed under different conditions. The data were obtained by modelingXPS spectra recorded at different depths of the passive film.

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NN passive films is 2.9 and 1.8 times that in the naturally formedpassive film, respectively. This indicates that the two electro-chemically formed films should be characterized by a higherdensity of metal vacancies than the naturally grown film, whichcould be the main reason for the observed difference in the pittingbehavior. Indeed, the Mott–Schottky measurements presentedin the following section of the paper will prove this assumption.

The Fe oxidation state depth profile analysis reveled thatthere are not any distinct differences in the distribution of Fe(II),Fe(III) and Fe(0) species in the three films. On the other hand,the results in Fig. 10 demonstrate that the two electrochemicallyformed passive films (ME-NN and ME-PB) show theaccumulation of Cr(VI) species in the outer part of the film.However, Cr(VI) was not detected in the naturally grownpassive film (Fig. 10, unmodified). This is quite understandableconsidering that the two electrochemically grown films areformed by the potentiodynamic polarization of the 316LVMsurface to high anodic potentials, where the oxidation of Cr(III)to Cr(VI) occurs (shoulder A2 in Fig. 1). Unlike Cr(VI), Fig. 10demonstrates that Cr(III) species are present in all three passivefilms, and that their corresponding depth distribution is verysimilar. Taking this into account, and also taking into accountthat there are not any distinct differences in the distribution ofFe species in the three films, and that the normalized oxygencontent in the electrochemically formed passive films is higherthan that in the naturally grown film (as discussed in theprevious paragraph), it seems that the formation of Cr(VI)species in the passive film is the origin of the observed dif-ference in pitting corrosion behavior presented in Figs. 2 and 4.Further, the fact that the thinner passive film formed in sodiumnitrate (NN) offered higher pitting resistance than the thickerfilm formed in phosphate buffer (PB), Fig. 2, indicates that thefilm thickness is not responsible for the difference in pittingcorrosion resistance between the two electrochemically formedfilms. However, we do not want to eliminate the possibility that

the film thickness could be partially responsible for the dif-ference in pitting resistance between the naturally and electro-chemically formed films.

Since possible incorporation of N- and P-species into thepassive layer can also contribute to the increased corrosionresistance of the investigated films, their presence in the formedpassive films was investigated by XPS. However, N was notdetected in the passive film formed in the sodium nitratesolution, while, on the other hand, P was detected in the filmformed in the phosphate buffer solution. Hence, the presence ofP in the ME-PB film could contribute to the observed increasein the film resistance (compared to the naturally grown film).Yet, the (N- and P-free) passive film formed in the sodiumnitrate solution offers the highest pitting corrosion resistance.This indicates that the presence of P in the ME-PB film makesonly a minor contribution to the increased corrosion resistanceof the film.

3.8. Semiconducting properties

It is known that passive films on SSs exhibit semiconductingbehavior [65,71–73], which significantly influences the result-ing corrosion properties of these materials. Relations betweensemiconducting properties and susceptibility of passive oxidefilms to localized corrosion has been well documented [74–76].Usually, electronic/semiconducting properties of these films areinvestigated using a capacitance technique [71,72,77]. Then, inthe interpretation of capacitance measurements, the Mott–Schottky approach is used, based on the assumption that thecapacitance response is controlled by the band bending. Thus,

Fig. 11. Mott–Schottky plots of 316LVM passive films. The graphs show theresponse of the unmodified 316LVM surface and the surfaces modified at 22 °Cin 0.1 M NaNO3 between −0.9 V and 0.9 V at a scan rate of 100 mV s−1 byapplying the specified number of potentiodynamic polarization sweeps. Themeasurements were done at 22 °C in 0.1 M phosphate buffer solution pH 7.4under the negative potential bias. Frequency=5 kHz, ac amplitude (rms):±10 mV. The data was corrected for the contribution of double layer capacitance.

104 A. Shahryari et al. / Materials Science and Engineering C 28 (2008) 94–106

the capacitive behavior of the passive film/electrolyte interfaceis assumed to be similar to that of a semiconductor/electrolyteinterface, and the total measured capacitance includes the spacecharge capacitance, CSC, and the Helmholtz double layercapacitance, CH in series: 1 /CT=1/CSC+1/CH. The Helmholtzdouble layer capacitance is considered to be constant [73].

The variation of space charge capacitance in the passive filmwith applied potential could be described by the well-knownMott–Schottky equation:

1

C2SC

¼ 2eeoeNdop

E � Efb � kTe

� �ð1Þ

where ɛ is the dielectric constant of the oxide film, ɛo is thevacuum permittivity (F cm−1), e is the elementary charge of anelectron (C), Ndop is the concentration of dopants (cm−3), Efb isthe flat-band potential (V), k is the Boltzmann constant (JK−1),and T is the temperature (K). If an oxide film behaves as asemiconductor that can be described by the Mott–Schottkymodel, the dependence between 1 /CSC

2 and E should give astraight line (positive for an n-type and negative for a p-typesemiconductor). An n-type behavior of SS passive films hasbeen correlated with the response of iron-oxides in the passivefilm, which are characterized by a non-stoichiometric compo-sition resulting in oxygen vacancies that contribute to the n-typesemiconductivity [74,78,79]. On the other hand, a p-typebehavior has been attributed to the response of chromium-oxides, and is characterized by the non-stoichiometry resultingin metal vacancies that contribute to the p-type semiconductivity[74,78,79].

The Mott–Schottky response of a naturally grown passivefilm (unmodified surface) and of passive films formed byapplying a specific number of polarization (modification) cycles(Fig. 11) shows the existence of several potential regionscharacterizing the semiconducting behavior of thecorresponding passive films. The existence of multiple Mott–Schottky regions has been explained by the existence of donor/acceptor states of different energy levels (deep and shallow)[71,78]. Fig. 11 shows that each passive film formed on the316LVM surface gives the same response in Regions I, II, IIIand V. These regions have been discussed elsewhere [80] and,thus, will not be discussed here. If we compare the Mott–Schottky (Fig. 11) and pitting (Fig. 2) curves of the unmodifiedsample, we will see that the pitting potential coincides well withthe n-type behavior in Region IV. Hence, it appears that on theunmodified surface (naturally formed passive film), oxygenvacancies act as pitting initiation sites. This is quite inagreement with the point defect theory [81,82], which statesthat the initial pitting reaction that occurs at the film/solutioninterface involves the adsorption of chloride ions into theseoxygen vacancies, followed by a Schottky-pair type reaction.This adsorption of chlorides leads to the generation of cationvacancies at the film/solution interface, followed by theirflux through the passive film to the film/metal interface. If theflux of these cation vacancies is so high that it cannot becompensated by the generation of cations at the metal/film

interface, cation vacancy condensates are formed. This, in turn,results in thinning of the passive film, or local detachment fromthe metal. Once the condensates have grown to a critical size,the film ruptures and rapid local pitting attack occurs.

Taking into account the point defect theory, our assumptionis that an increase in pitting corrosion resistance of the materialcould be achieved by increasing the concentration of metalvacancies in the passive film. On the Mott–Schottky plot(Fig. 11), this change would be manifested as a transition froman n- to a p-type semiconductivity in the pitting susceptibleRegion IV. Indeed, Fig. 11 shows that by modifyingthe 316LVM surface by cyclic potentiodynamic polarization,the n-type semiconductivity in Region IV gradually transformsinto the p-type semiconductivity, finally resulting in mergerof Regions IVand V (Fig. 11, diamond symbols). Consequently,the resulting pitting potential also gradually increases, asdemonstrated in Fig. 3. The gradual transition from the n-typeto the p-type semiconductivity is related to the gradualenrichment of the passive film with metal vacancies. Thisprocess occurs by the cyclic potentiodynamic formation ofthe passive film under the conditions presented in Fig. 1.Namely, by polarizing the 316LVM surface at high anodicpotentials, Cr(VI) species are formed. With cyclization,progressively more of these species remain ‘arrested’ in thegrowing passive film (inset to Fig. 1 and Fig. 10). On the otherhand, only Cr(III), but not Cr(VI), species are formed duringthe natural growth of the film (unmodified sample), which isclearly demonstrated in Fig. 10. The formation of Cr(VI)species in the form of Cr2O7

2−increases the oxygen content inthe passive film (Fig. 9b), producing an increase in the densityof metal vacancies, VM

3−, according to [83]:

Cr IIIð ÞYCr VIð Þ þ V 3�M þ 3e� ð2Þ

105A. Shahryari et al. / Materials Science and Engineering C 28 (2008) 94–106

This, in turn, leads to the change in the Mott–Schottkybehavior in Region IV (Fig. 11), and finally to the increasedpitting resistance of the passive film (Fig. 2).

4. Conclusion

It was demonstrated that the modification of the biomedicalgrade 316LVM stainless steel surface by cyclic potentiody-namic polarization in sodium nitrate or phosphate buffer resultsin the formation of passive oxide films that offer a significantlyhigher pitting corrosion resistance than the naturally grownpassive film. The film formed in the nitrate electrolyte showedto be completely resistant to pitting in physiological simulatingelectrolytes (PBS and Hank's) even at high temperatures andafter sample sterilization, and to offer very high pittingcorrosion resistance in the pure saline electrolyte, even atelevated chloride concentrations. It is also worth to mention thatthe film formed in the nitrate electrolyte maintained the same(high) pitting resistance even after two months of constantexposure to the pure saline solution.

The capacitance analysis demonstrated that the majordifference between the electrochemically formed and naturallygrown passive film is in the type of semiconductivity in thepotential region where pitting on the unmodified surface occurs.The XPS measurements showed that this is due to the presenceof electrochemically formed Cr(VI)-species in the outer part ofthe electrochemically formed passive film. Namely, theelectrochemical formation of the passive film results in adecrease in the density of oxygen vacancies, which act as pittinginitiation sites, and their ‘replacement’ by metal vacanciesformed by the oxidation of Cr(III) to Cr(VI). In thisconfiguration the outer Cr(VI)-rich oxide layer behaves ascation selective which, in turn, results in the increased pittingcorrosion resistance of the corresponding passive film.

Acknowledgements

Grateful acknowledgment is made to the Natural Science andEngineering Research Council of Canada and the McGill Centrefor Biorecognition and Biosensors for support of this research.

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